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Fabrication of W–Cu composite by shock consolidation of Cu-coated W powders
Journal of Alloys and Compounds 657 (2016) 215e223
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Fabrication of WeCu composite by shock consolidation of Cu-coated W powders Qiang Zhou, Pengwan Chen* State Key Laboratory of Explosion Science and Technology, Beijing Institute of Technology, Beijing 100081, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 3 August 2015 Received in revised form 3 October 2015 Accepted 6 October 2015 Available online 14 October 2015
Nearly fully dense (>98% theoretical density) WeCu composites, using Cu-coated W powders, have been fabricated by shock consolidation. The composites are free of cracks, and shows homogenous distribution of W and a Cu network structure. The high density is obtained under relatively low shock pressure of ~1 GPa without preheating; because the sintering is only related to CueCu sintering. The thermal (thermal conductivity) and mechanical (quasi-static compression and hardness) properties of the composites are measured experimentally. It is found that the bonding of W and Cu is determined during the electroless coating process and increasing in shock intensity doesn't improve the bonding strength. Inappropriate coating process and subsequent oxidation leads to the weak bonding and inferior thermal performance in this work. The Cu network as a continuous phase significantly influences the mechanical response and the weak bonding strength between W and Cu leads to axial splitting as a primary failure mode. After increasing the shock intensity, nano-sized W particles are formed at the center due to the particle crushing and jet penetration. The fragmentation of W particles decreases the hardness at the center because good bonding is formed due to the high surface energy of nano-sized particle. © 2015 Elsevier B.V. All rights reserved.
Keywords: Cu-coated W powder Shock consolidation WeCu composites
1. Introduction Tungsten-copper (WeCu) composites combine the excellent properties of both components, such as low thermal expansion, high melting temperature, and high temperature strength of W, high thermal and electrical conductivity of Cu [1]. Such properties make WeCu composites be widely used as heat-sink material, ultra-high voltage electric contact material and warhead materials [2]. Due to the high differences in densities and melting points, as well as the poor mutual solubility, WeCu composites are typical materials fabricated by powder metallurgy (PM) technology. For the fabrication of pseudo-alloys such as WeCu, high sintering temperature and long holding time are necessary to improve the densification. But grain coarsening cannot be avoided and Cu may leach out from the skeleton that leads to Cu segregation and results in non-homogenous microstructure and poor product performance [3]. Several novel techniques have been explored to overcome the deterioration due to the high sintering temperature and long
* Corresponding author. E-mail address: [email protected] (P. Chen). http://dx.doi.org/10.1016/j.jallcom.2015.10.057 0925-8388/© 2015 Elsevier B.V. All rights reserved.
holding time, which includes microwave sintering [3], resistance sintering [4,5] and cold rolling [6]. However, the sintering temperature and holding time are still substantial. Hot-shock consolidation technique has been successfully utilized to fabricate pure fine-grained W [7] and WeCu composites [8] at a relatively low temperature. It was found that the homogenous distribution of Cu is the key factor for the consolidation of WeCu powder, which is dominated by preheating temperature. High temperature improves the distribution of Cu and enhances the bonding of WeCu, but also leads to agglomerating of Cu while exceeding a specific value. It is suggestive that homogenous distribution of W and Cu may lead to better sintering, which is in accordance with the opinion of Xu et al. [9]. Many methods have been explored to obtain homogenous distribution of W in Cu matrix, such as mechanical alloying, mechano-chemical processing [10] and thermo-mechanical method [11]. However, ball milling in these methods often takes long time; and impurities could be introduced during the milling process, which will decrease the performance of WeCu composites. Some novel methods, such as wet-chemical methods [12] and spray drying [13] have also been tried to synthesize WeCu nanopowders, in which the conventional mechanical milling is replaced by wet mixing. However, a series of compound salts may form in an aqueous solution, and the
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composition of the precursors as well as the resultant WeCu powders may depend greatly on the stoichiometric amount in the solution. Cu can be uniformly and densely coated on the surface of W powder particles utilizing chemical method, and this Cu-coated W powders can be sintered at a relatively low temperature with a short sintering duration. Zhang et al. [14] fabricated high-density W-20 wt.%Cu by hot-press sintering at 950 C-100MPa-2 h using Cu-coated W powders. It was claimed that a network structure of high-purity Cu and a homogenous distribution of W particles are key factors for excellent thermal performance of WeCu composites. Ibrahim et al. [1] compared the sintering of Cu-coated W powders to admixed WeCu powders. The results show that the coated powders have greater compressibility and higher performance than the uncoated ones. High sintering temperature and long duration are still inevitable while using conventional PM methods to sinter Cu-coated W powders. However, shock consolidation seems to be a promising way because powders experience a pressures level exceeding 1 GPa at extremely high strain rates (107e108 s 1) during shock consolidation; and such a rapid deformation induces particle surface heating and can result in extremely high cooling rates of 109 K/s near the surface [15], which allows the possibility of surface melting while the interior of particle remains relatively cool [16]. In this paper, shock consolidation was applied to sinter the Cu-coated W powders without preheating. The microstructure, thermal performance and mechanical properties of WeCu composites were investigated. 2. Experimental procedure The Cu-coated W powders (W-20 wt.% Cu) were prepared by chemical reduction (electroless coating). The coating process was carried out using a solution consists of copper sulfate
(CuSO4$5H2O), complexing agent (KNaC4O6H4$4H2O) and stabilizing agent (C10H14N2Na2O8$2H2O), which was made in deionized water using alkaline solution to adjust the pH value. The mixed solution was stirred at a temperature of 70 C, and the methanol with a concentration of 30 ml/L and W powders were poured into the solution simultaneously; then the Cu was deposited on W powders. The precipitated WeCu powders were washed with deionized water and dried under vacuum. The distribution of particle size of Cu-coated W powders is uniform and with an average size of ~20 mm, as shown in the SEM images in Fig. 1. As shown in Fig. 1(b), the W particles are coated with a layer of Cu with some air holes in it. As can been seen in Fig. 1(c), the Cu layer is not quite uniform, with the thickness in the range from 1.5 mm to 5 mm. The high depositing rate due to high stirring temperature is responsible for the coarse structure of Cu layer. The magnified view in Fig. 1(d) shows the gap at the interface between W and Cu, indicating a weak bonding. The EDS analysis of Cu-coated W powders shows weak W peak, as illustrated in Fig. 2, indicating W particle is not completely wrapped or the Cu layer is thin, which agrees with Fig. 1(c). And some impurities such as O and S were also observed. The oxide on copper surface is due to oxidation during the storage and packing stage; and the sulfured is unreliable as its intensity is within the noise. The powders were compacted into 18 mm (inner) diameter steel tubes 100 mm in length, with steel end plug welded to both ends. These tubes were placed in PVC cylinders with various diameters, and filled with explosives as illustrated schematically in Fig. 3. Cylindrical shock-wave consolidation arrangements characteristic of that shown in Fig. 3 have been described in detail by Prümmer [17] and Meyers and Wang [18]. The powder explosive expanded ammonium nitrate with a detonation velocity of 2300 m/s and a liquid explosive nitromethane with a detonation velocity of 6300 m/s, were used in these experiments. The shock pressures
Fig. 1. Particle size distribution of investigated Cu-coated W powders (a); magnified view of powder particles (b); cross-section of Cu-coated W powders (c); and the gap at the interface between W and Cu.
Q. Zhou, P. Chen / Journal of Alloys and Compounds 657 (2016) 215e223
Fig. 2. EDS analysis of the Cu-coated W powders.
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associated with these explosive consolidation experiments were ~1 GPa and ~12 GPa for the powder and liquid explosive, respectively. Fig. 4(a) shows a shock consolidation assembly placed on sand pile; Fig. 4(b) is a snapshot of powder containers sealed by steel plugs welded to both ends. Fig. 4(c) shows the recovered sample containers (with reduced diameter) after shock loading. Fig. 4(d) shows the specimens machined out from the compacted containers. The densities of the sintered samples were measured by Archimedes' method. Scanning electron microscopy (SEM) with an energy dispersive spectrum (EDS) and optical microscopy were used to observe the morphology of the fractured surface and the microstructure. The samples for optical and SEM observation were ground and polished to mirror, then etched with a solution of
Fig. 3. Schematic of the experimental setup for shock consolidation and the detonation process.
Fig. 4. Shock consolidation assembly placed in a sand pit (a); the powder containers sealed with plugs welded to both ends (b); as recovered compacted containers with reduced diameter (c); and the specimen machined out from container (d).
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10 g K3Fe(CN)6, 10 g KOH and 100 ml water. X-ray diffraction with Cu Ka radiation (40 kV, 40 mA) was used for phase analysis. The thermal diffusivity was measured by the laser flash method on a LFA447 at room temperature, using a cylindrical sample with a diameter of 10 mm and a thickness of 2 mm. The quasi-static compression tests were carried out using cylindrical samples with a diameter of 10 mm and a height of 10 mm at a strain rate of 10 3 s 1. The hardness measurements were performed at a load of 5 kg for 30 s using a Vickers hardness tester. 3. Results and discussion 3.1. XRD analysis
Fig. 5. XRD patterns of the Cu-coated W powder and the sintered samples.
Fig. 5 shows the XRD analysis of initial powder and the sintered samples, and no impurities are observed. As shown in the inset of Fig. 5, the intensity of W peak for AN-7 is higher than the one for powder. This is because the coreeshell structure of the Cu coated W powder. But the intensity of W peaks for NM-1 is weaker than AN-7. This is because submicron particles were formed at the central region under extreme high pressure, leading to reduction of
Fig. 6. Light microscopy image of WeCu composites AN-7 (a); magnified views of the microstructure at the center (b) and the peripheral region (c) of AN-7. Light microscopy image of WeCu composites NM-1 (d); magnified views of the microstructure at the center (e) and the peripheral region (f) of NM-1.
Fig. 7. A typical image of pit due to weak bonding between W and Cu (a); magnified view of nano-sized particle formed at the center of NM-1 (b).
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Table 1 The properties of WeCu composites sintered with various conditions. No.
Shock pressure (GPa)
Sintered density %
Yield stress (MPa)
Thermal conductivity (W/mK)
AN-7 NM-1
1 12
98.9 98.4
554.52 504.64
144.82 145.62
intensity. The reason for the formation of nano-sized particles will be discussed in the section below. The W peaks from XRD profiles for initial powder, AN-7 and NM-1 were fitted using Pseudo-Voigt Function to estimate the peak widths. The results show no obvious peak broadening for the sintered samples. That implies that the submicron particles in NM-1 are formed by mechanical fragmentation, not grain refinement. 3.2. Microstructure analysis Fig. 6(a) and Fig. 6(d) are light microscopy images of AN-7 and NM-1, respectively. Fig. 6(c) and (f) show uniform distribution of W particles in Cu matrix, where no aggregations of W were observed. There are numerous dark patches distributed across the cross section of AN-7, and at the peripheral region of NM-1 as well. The samples are essentially free of porosity with a measured density > 98%T.D. (theoretical density). The magnified view by SEM shows the dark patch indicated by arrow in Fig. 7(a) is not a pore but pit. The formation of pores during shock consolidation is quite different from conventional powder metallurgy method. For shock consolidation, pores are eliminated through plastic deformation and particle fracture which are driven by shock wave instantaneously. So the pores in WeCu powders tend to retain between W particles due to insufficient deformation of W, which has been verified in hot-shock consolidation of WeCu [8]. But in this work, the WeW contact is avoided by using Cu-coated W powder. Furthermore, the size of dark patch is too large to be a pore, especially under the high pressure and high strain rate condition. However, it does exist pores due to the shrink of melted Cu under extremely high cooling rates, as indicated by the circle in Fig. 7(a). The angular pit is the evidence where a W particle existed, indicating a weak bonding between W and Cu. It is believed the weak bonding mainly comes from coating process, and the oxidation during the storage and packing stage makes it worse. The high
Fig. 8. Magnified view of center region of NM-1 shows various rupture modes: (a) fragments due to jetting penetration; (b) fragments due to impact between W particles; and (c) crushed zone at the surface.
depositing rate of Cu due to high stirring temperature lead to coarse Cu layer during the coating process, as shown in Fig. 1(b). The coarse Cu layer and the low surface energy of W particle due to its relatively large size are responsible for the weak bonding between Cu layer and W particle. The densification mechanism of Cu-coated W powders is different from the one of hot-shock consolidation with mechanically mixed WeCu powder, which is dominated by densification of W particles [8]. The densification of W particles is dominated by shock pressure; it means high-pressure is needed to promote void collapse and plastic deformation of W particles. For Cu-coated W powders, the sintering is only related to CueCu sintering, which can be easily achieved under relatively low shock pressure without preheating. This indicates that the bonding of W and Cu, which is crucial to the performance of composite, is determined by the electroless coating process. As shown in Table 1, high relative density (>98%) is obtained with a shock intensity of ~1 GPa. However, the sintering conditions required for hot-shock consolidation method [8] are high pressure (>3 GPa) and high preheating temperature (900e1000 C). By using the Cu-coated W powders, sintering mode is changed and there is no requirement of preheating. Therefore, the shock pressure is reduced and the density of the composites is improved. On cylindrical shock consolidation shock waves obviously converge and the pressures grow accordingly toward the powder center. But, during passage of the shock wave, the powders become deformed and get sintered. Energy is consumed and the shock wave should attenuate while advancing. A combination of the two effects leads to various shock front shapes, which determines the uniformity of compaction. The ideal case is when the two opposing effects counterbalance, with a hollow cone-shaped shock front as shown in Fig. 3. In such a case a uniform pressure distributes over the crosssection and a uniform consolidation is obtained. The AN-7 shows a uniform microstructure across the cross section, as shown in Fig. 6 (a), (b) and (c), indicating an ideal condition. If the amount of explosive or its detonation velocity is too high, as the case of NM-1, the Mach-stem will be formed. Under such conditions, the pressure increases continuously towards the center, cracking and melting may occur due to the extremely high pressure to tens or hundreds of GPa. This phenomenon has been observed during consolidation of metallic powders and has been studied in Refs. [19,20]. At the peripheral region for both AN-7 and NM-1, the morphologies are quite similar with the W particle in the size identical to initial powders, comparing Fig. 6(c) with Fig. 6(f). At the center region, NM-1 shows a fully densified structure without any pits or pores, as shown in Fig. 6(e). Furthermore, plenty of ultra-fined W particles are observed in Fig. 6(e), the size of which is much smaller than initial powder and is identified as nanoscale by Fig. 7(b). It is evident, from the magnified view in Fig. 8, that these nanosized W particles were formed due to the particle crushing and jet penetration. The schematic diagram in Fig. 9 shows the formation of nano-sized W particle at the center region in NM-1. While two Cu-coated W particles impact each other at high velocity driven by extremely high pressure, a Cu jet will occur due to the oblique impact of Cu layers. The ejected jet penetrates adjacent W particles at extremely high velocity, and particle crushing occurs, rather than deformation or melting due to the brittleness and high melting
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Fig. 10. Stressestrain curves from samples of quasi-static compression test.
observed in Mach-stem for metallic powders, have been avoided due to ductility of Cu and high-melting point of W, and especially due to using coating method. 3.3. Thermal and mechanical properties of WeCu composites As shown in Table 1, the thermal conductivities are 144.82 W/ mK and 145.62 W/mK for AN-7 and NM-1, respectively, which are lower than the theoretical value (237 W/mK). A continuous network of Cu, high sintering density and good interfacial contact between W and Cu are the keys to excellent thermal properties of WeCu composites [14]. As mentioned above, the weak bonding between W and Cu should be responsible for the poor thermal performance of the composites in this work. The quasi-static compression test results, shown in Fig. 10, reveal the slight difference between the two samples. These cylindrical samples have dimensions of 10 mm (length) 10 mm (diameter). No extensometer was used and the results were corrected from loadedisplacement plots by subtracting machine effects. The yield stresses are listed in Table 1. For the explosively consolidated powder mixture compact, it is suggested that the yield strength is determined by the continuous phase [21]. In this Fig. 9. Schematic diagram of the formation of nano W particles.
point of W. The small W particles are crushed into pieces while the crushing for a big particle is limited to local region. The crushed W fragments burst out and disperse into Cu matrix. This is the primary rupture mode. Simultaneously, the impact between W particles also leads to particle crushing. The crushed zone due to the normal impact is restrained to surface and the fragments due to oblique impact are propelled into Cu matrix. But these two rupture modes are limited to local region, because most kinetic energy has been deposited as deformation and jetting of Cu. The corresponding rupture modes illustrated in Fig. 9 are indicated in Fig. 8. As indicated by the arrow (a) in Fig. 8, a concave on a W particle is filled with fragments without contact with any other particles around; it is obviously a proof of jet penetration. As indicated by the arrow (b), the splashed fragments into Cu matrix show an oblique impact between W particles. And the arrow (c) in Fig. 8 shows a cursed zone due to the normal impact, with the fragments staying between W particles. The cracking and melting, which are usually
Fig. 11. Variation of hardness of the investigated samples from center to periphery.
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Fig. 12. Fracture morphology of WeCu composites after quasi-static compression test: (a) SEM image of fractured NM-1; (b) SEM image of fractured AN-7; (c) SEM image showing smooth surface of Cu due to sliding of W and (d) backscattered electron micrograph showing a W particle pulled out from the Cu matrix (dark gray for Cu and light gray for W).
work, the samples consolidated with different pressure have similar yield stress, ~500 MPa. However, the yield strength of the compact with higher shock pressure is ~50 MPa lower than the one with lower pressure, but its ductility is slightly improved. This will be discussed further below. Fig. 11 shows the variation of the hardness across the crosssection of the sample. It can be seen that the hardness of AN-7 remains consistent throughout the consolidate indicating that densification was quite uniform. But for NM-1, the graph shows a decrease from periphery towards center and a sharp drop to 107 Hv at the center. In general, the hardness increases with decreasing particle size for particulate-reinforced metal matrix composites. But the reverse behavior also has been reported [22], where the deterioration is interpreted by the inhomogeneous distribution of submicron particle. For the WeCu composites fabricated with Cucoated W powder, no W skeleton is formed during the sintering. And the indentation deformation is subjected to Cu matrix because the hardness of W particles is much higher than Cu matrix. Only if the Cu between W particles is crowded out, the W particle will be involved in resisting to the indentation. At the periphery of NM-1 where the W particles are micron-sized, Cu is crowded out during the indentation due to the weak bonding between W and Cu. And the hardness accordingly increases due to the participation of W particle. But at the center where part of the W particles are nanosized, high bonding strength is formed due to the high surface energy of nano-particles. The high bonding strength, as well as the increased WeCu interfaces, absorbs the mechanical energy that will lead to WeCu debonding and Cu extrusion; then hardness is close to the one of Cu matrix due to the lack of W participation. The high bonding strength and increasing WeCu interfaces as a result of particle fragmentation is also the reason for the improvement in bulk's ductility. The fracture type of both samples can be distinctly concluded: axial splitting fracture parallel to the loading direction. The
fractured samples were observed by scanning electron microscopy, as the general fractographs of NM-1 and AN-7 shown in Fig. 12(a) and Fig. 12(b), respectively. This failure mode is similar to the one of explosively consolidated Ni/WeAl composites under dynamic compression [21]. The discontinuously distributed W particles dispersed in the continuous Cu matrix can be considered as rigid during plastic deformation in uniaxial compression loading because its hardness significantly exceeds that of Cu. The separation is initiated at the interface between the W particle and Cu matrix, which is shown in the elemental mapping result of Fig. 13. During compression loading, the W particles move and cluster in the Cu matrix, verified by the scratch shown in Fig. 12(c). As vertical compression continues, the W particles are crowded out and move laterally, as indicated by the arrow in Fig. 12(d). Then the cracks parallel to loading direction are created due to low bonding strength between W and Cu, and propagate through the entire sample as axial splitting. The two samples consolidated with different pressure show the similar fracture mode, indicating shock pressure has no help to improve WeCu bonding strength. 4. Conclusions The following are the principal conclusions drawn from the experimental work performed on shock consolidation of Cu-coated W powder: 1. The Cu-coated W powders are prepared by chemical reduction (electroless coating). The high depositing rate of Cu during coating process and the relatively large size of W particle are responsible for the weak bonding in this work. The impurities are introduced by oxidation during the storage and packing, which also decrease the performance of WeCu composites. 2. The Cu-coated W powders were consolidated to a high relative density (exceeding 98%) by the shock wave without preheating.
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Fig. 13. Elemental mapping of the fracture surface: (a) backscattered electron micrograph of the region with a crack; (b) elemental distribution of Cu; (c) elemental distribution of W and (d) combined elemental distribution of W and Cu.
The sintering is only related to CueCu sintering due to the homogenous distribution of Cu on the W surface; and no W skeleton is formed during the sintering. Then the strength of consolidated composite is primarily determined by the Cu phase, exhibiting a lower value. The weak WeCu bonding from coating process, as well as the impurities, leads to axial splitting as a primary failure mode and inferior thermal performance. 3. The shock waves with the intensities of ~1 GPa and ~12 GPa were applied to consolidate Cu-coated W powders, respectively. The results show no difference in the bulk's performance but big difference in microstructure. The composite fabricated under 12 GPa shows an inhomogeneous microstructure along the cross section; and nano-sized W particles are formed at the center due to the particle crushing and jet penetration, as a result of Mach effect. The XRD result shows that the nano-sized particle is the result of mechanical fragmentation not grain refinement. The fragmentation of W particles decreases the hardness at the center because good bonding is formed due to the high surface energy of nano-sized particle. Acknowledgments The authors would like to express their thanks for the financial support of National Natural Science Foundation of China under grant nos. 11172043, 11402026 and 11221202. And the authors also express the appreciation to Professor C.W. Tan from school of Material Science and Technology in Beijing Institute of Technology, for the support on powder preparation. References [1] A. Ibrahim, M. Abdallah, S.F. Mostafa, A.A. Hegazy, An experimental investigation on the WeCu composites, Mater. Des. 30 (2009) 1398e1403. [2] Z.J. Zhou, J. Du, S.X. Song, Z.H. Zhong, C.C. Ge, J. Alloys Comp. 428 (2007) 146e150.
[3] A. Modal, A. Upadhyaya, D. Agrawal, Comparative study of densification and microstructural development in We18Cu composites using microwave and conventional heating, Mater. Res. Innov. 14 (5) (2010) 355e360. [4] Z.J. Zhou, Y.S. Kwon, Fabrication of WeCu composites by resistance sintering under ultra-high pressure, J. Mater. Process. Technol. 168 (2005) 107e111. [5] Z.J. Zhou, J. Du, S.X. Song, Z.H. Zhong, C.C. Ge, Microstructural characterization of W/Cu functionally graded materials produced by a one-step resistance sintering method, J. Alloys Compd. 428 (2007) 146e150. [6] N. Yang, Z. Wang, L. Chen, Y. Wang, Y.B. Zhu, A new process for fabricating W15 wt.% Cu sheet by sintering, cold rolling and re-sintering, Int. J. Refract. Metals Hard Mater. 28 (2010) 198e200. [7] Q. Zhou, P. Chen, Fabrication and characterization of pure tungsten using the hot-shock consolidation, Int. J. Refract. Metals Hard Mater. 42 (2014) 215e220. [8] Q. Zhou, P. Chen, Characterization of fine-grained W-10 wt% Cu composite fabricated by hot-shock consolidation, Int. J. Refract. Metals Hard Mater. 52 (2015) 137e142. [9] L. Xu, M. Yan, Y. Xia, J. Peng, W. Li, L. Zhang, C. Liu, G. Chen, Y. Li, Influence of copper content on the property of CueW alloy prepared by microwave vacuum infiltration sintering, J. Alloys Compd. 592 (2014) 202e206. [10] J. Cheng, P. Song, Y. Gong, Y. Cai, Y. Xia, Fabrication and characterization of We15Cu composite powders by a novel mechano-chemical process, Mater. Sci. Eng. A 488 (2008) 453e457. [11] S. Hong, B. Kim, Fabrication of W-20 wt% Cu composite nanopowder and sintered alloy with high thermal conductivity, Mater. Lett. 57 (18) (2003) 2761e2767. [12] L. Wan, J. Cheng, P. Song, Y. Wang, T. Zhu, Synthesis and characterization of WeCu nanopowders by a wet-chemical method, Int. J. Refract. Metals Hard Mater. 29 (4) (2011) 429e434. [13] X. Shi, H. Yang, S. Wang, G. Shao, X. Duan, Z. Xiong, T. Wang, Characterization of We20Cu ultrafine composite powder prepared by spray drying and calcining-continuous reduction technology, Mater. Chem. Phys. 104 (2007) 235e239. [14] Lianmeng Zhang, Wenshu Chen, Guoqiang Luo, Pingan Chen, Qiang Shen, Chuanbin Wang, Low-temperature densification and excellent thermal properties of WeCu thermal-management composites prepared from coppercoated tungsten powders, J. Alloys Compd. 588 (2014) 49e52. [15] W.F. Gourdin, Energy deposition and microstructural modification in dynamically consolidated metal powders, J. Appl. Phys. 55 (1984) 172. [16] T.G. Nieh, P. Luo, W. Nellis, D. Lesuer, D. Benson, Dynamic compaction of aluminum nanocrystals, Acta Mater. 44 (9) (1996) 3781e3788. [17] R. Prümmer, Explosivverdichten Pulvriger Substanzen, Springer-Verlag, Berlin, 1987. [18] M.A. Meyers, S.L. Wang, An improved method for shock consolidation of
Q. Zhou, P. Chen / Journal of Alloys and Compounds 657 (2016) 215e223 powders, Acta Metall. 36 (1988) 925e936. [19] R.A. Pruemmer, T. Balakrishna Bhat, K. Siva Kumar, K. Hokamoto, Explosive Compaction of Powders and Composites, New Hampshire: Science Publishers, 2006. [20] S.S. Batsanov, Effects of Explosive on Materials-modification and Synthesis under High-pressure Shock Compression, Springer-Verlag, New York, 1994. [21] C.T. Wei, E. Vitali, F. Jiang, S.W. Du, D.J. Benson, K.S. Vecchio, N.N. Thadhani,
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Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Fabrication of WeCu composite by shock consolidation of Cu-coated W powders Qiang Zhou, Pengwan Chen* State Key Laboratory of Explosion Science and Technology, Beijing Institute of Technology, Beijing 100081, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 3 August 2015 Received in revised form 3 October 2015 Accepted 6 October 2015 Available online 14 October 2015
Nearly fully dense (>98% theoretical density) WeCu composites, using Cu-coated W powders, have been fabricated by shock consolidation. The composites are free of cracks, and shows homogenous distribution of W and a Cu network structure. The high density is obtained under relatively low shock pressure of ~1 GPa without preheating; because the sintering is only related to CueCu sintering. The thermal (thermal conductivity) and mechanical (quasi-static compression and hardness) properties of the composites are measured experimentally. It is found that the bonding of W and Cu is determined during the electroless coating process and increasing in shock intensity doesn't improve the bonding strength. Inappropriate coating process and subsequent oxidation leads to the weak bonding and inferior thermal performance in this work. The Cu network as a continuous phase significantly influences the mechanical response and the weak bonding strength between W and Cu leads to axial splitting as a primary failure mode. After increasing the shock intensity, nano-sized W particles are formed at the center due to the particle crushing and jet penetration. The fragmentation of W particles decreases the hardness at the center because good bonding is formed due to the high surface energy of nano-sized particle. © 2015 Elsevier B.V. All rights reserved.
Keywords: Cu-coated W powder Shock consolidation WeCu composites
1. Introduction Tungsten-copper (WeCu) composites combine the excellent properties of both components, such as low thermal expansion, high melting temperature, and high temperature strength of W, high thermal and electrical conductivity of Cu [1]. Such properties make WeCu composites be widely used as heat-sink material, ultra-high voltage electric contact material and warhead materials [2]. Due to the high differences in densities and melting points, as well as the poor mutual solubility, WeCu composites are typical materials fabricated by powder metallurgy (PM) technology. For the fabrication of pseudo-alloys such as WeCu, high sintering temperature and long holding time are necessary to improve the densification. But grain coarsening cannot be avoided and Cu may leach out from the skeleton that leads to Cu segregation and results in non-homogenous microstructure and poor product performance [3]. Several novel techniques have been explored to overcome the deterioration due to the high sintering temperature and long
* Corresponding author. E-mail address: [email protected] (P. Chen). http://dx.doi.org/10.1016/j.jallcom.2015.10.057 0925-8388/© 2015 Elsevier B.V. All rights reserved.
holding time, which includes microwave sintering [3], resistance sintering [4,5] and cold rolling [6]. However, the sintering temperature and holding time are still substantial. Hot-shock consolidation technique has been successfully utilized to fabricate pure fine-grained W [7] and WeCu composites [8] at a relatively low temperature. It was found that the homogenous distribution of Cu is the key factor for the consolidation of WeCu powder, which is dominated by preheating temperature. High temperature improves the distribution of Cu and enhances the bonding of WeCu, but also leads to agglomerating of Cu while exceeding a specific value. It is suggestive that homogenous distribution of W and Cu may lead to better sintering, which is in accordance with the opinion of Xu et al. [9]. Many methods have been explored to obtain homogenous distribution of W in Cu matrix, such as mechanical alloying, mechano-chemical processing [10] and thermo-mechanical method [11]. However, ball milling in these methods often takes long time; and impurities could be introduced during the milling process, which will decrease the performance of WeCu composites. Some novel methods, such as wet-chemical methods [12] and spray drying [13] have also been tried to synthesize WeCu nanopowders, in which the conventional mechanical milling is replaced by wet mixing. However, a series of compound salts may form in an aqueous solution, and the
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composition of the precursors as well as the resultant WeCu powders may depend greatly on the stoichiometric amount in the solution. Cu can be uniformly and densely coated on the surface of W powder particles utilizing chemical method, and this Cu-coated W powders can be sintered at a relatively low temperature with a short sintering duration. Zhang et al. [14] fabricated high-density W-20 wt.%Cu by hot-press sintering at 950 C-100MPa-2 h using Cu-coated W powders. It was claimed that a network structure of high-purity Cu and a homogenous distribution of W particles are key factors for excellent thermal performance of WeCu composites. Ibrahim et al. [1] compared the sintering of Cu-coated W powders to admixed WeCu powders. The results show that the coated powders have greater compressibility and higher performance than the uncoated ones. High sintering temperature and long duration are still inevitable while using conventional PM methods to sinter Cu-coated W powders. However, shock consolidation seems to be a promising way because powders experience a pressures level exceeding 1 GPa at extremely high strain rates (107e108 s 1) during shock consolidation; and such a rapid deformation induces particle surface heating and can result in extremely high cooling rates of 109 K/s near the surface [15], which allows the possibility of surface melting while the interior of particle remains relatively cool [16]. In this paper, shock consolidation was applied to sinter the Cu-coated W powders without preheating. The microstructure, thermal performance and mechanical properties of WeCu composites were investigated. 2. Experimental procedure The Cu-coated W powders (W-20 wt.% Cu) were prepared by chemical reduction (electroless coating). The coating process was carried out using a solution consists of copper sulfate
(CuSO4$5H2O), complexing agent (KNaC4O6H4$4H2O) and stabilizing agent (C10H14N2Na2O8$2H2O), which was made in deionized water using alkaline solution to adjust the pH value. The mixed solution was stirred at a temperature of 70 C, and the methanol with a concentration of 30 ml/L and W powders were poured into the solution simultaneously; then the Cu was deposited on W powders. The precipitated WeCu powders were washed with deionized water and dried under vacuum. The distribution of particle size of Cu-coated W powders is uniform and with an average size of ~20 mm, as shown in the SEM images in Fig. 1. As shown in Fig. 1(b), the W particles are coated with a layer of Cu with some air holes in it. As can been seen in Fig. 1(c), the Cu layer is not quite uniform, with the thickness in the range from 1.5 mm to 5 mm. The high depositing rate due to high stirring temperature is responsible for the coarse structure of Cu layer. The magnified view in Fig. 1(d) shows the gap at the interface between W and Cu, indicating a weak bonding. The EDS analysis of Cu-coated W powders shows weak W peak, as illustrated in Fig. 2, indicating W particle is not completely wrapped or the Cu layer is thin, which agrees with Fig. 1(c). And some impurities such as O and S were also observed. The oxide on copper surface is due to oxidation during the storage and packing stage; and the sulfured is unreliable as its intensity is within the noise. The powders were compacted into 18 mm (inner) diameter steel tubes 100 mm in length, with steel end plug welded to both ends. These tubes were placed in PVC cylinders with various diameters, and filled with explosives as illustrated schematically in Fig. 3. Cylindrical shock-wave consolidation arrangements characteristic of that shown in Fig. 3 have been described in detail by Prümmer [17] and Meyers and Wang [18]. The powder explosive expanded ammonium nitrate with a detonation velocity of 2300 m/s and a liquid explosive nitromethane with a detonation velocity of 6300 m/s, were used in these experiments. The shock pressures
Fig. 1. Particle size distribution of investigated Cu-coated W powders (a); magnified view of powder particles (b); cross-section of Cu-coated W powders (c); and the gap at the interface between W and Cu.
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Fig. 2. EDS analysis of the Cu-coated W powders.
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associated with these explosive consolidation experiments were ~1 GPa and ~12 GPa for the powder and liquid explosive, respectively. Fig. 4(a) shows a shock consolidation assembly placed on sand pile; Fig. 4(b) is a snapshot of powder containers sealed by steel plugs welded to both ends. Fig. 4(c) shows the recovered sample containers (with reduced diameter) after shock loading. Fig. 4(d) shows the specimens machined out from the compacted containers. The densities of the sintered samples were measured by Archimedes' method. Scanning electron microscopy (SEM) with an energy dispersive spectrum (EDS) and optical microscopy were used to observe the morphology of the fractured surface and the microstructure. The samples for optical and SEM observation were ground and polished to mirror, then etched with a solution of
Fig. 3. Schematic of the experimental setup for shock consolidation and the detonation process.
Fig. 4. Shock consolidation assembly placed in a sand pit (a); the powder containers sealed with plugs welded to both ends (b); as recovered compacted containers with reduced diameter (c); and the specimen machined out from container (d).
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10 g K3Fe(CN)6, 10 g KOH and 100 ml water. X-ray diffraction with Cu Ka radiation (40 kV, 40 mA) was used for phase analysis. The thermal diffusivity was measured by the laser flash method on a LFA447 at room temperature, using a cylindrical sample with a diameter of 10 mm and a thickness of 2 mm. The quasi-static compression tests were carried out using cylindrical samples with a diameter of 10 mm and a height of 10 mm at a strain rate of 10 3 s 1. The hardness measurements were performed at a load of 5 kg for 30 s using a Vickers hardness tester. 3. Results and discussion 3.1. XRD analysis
Fig. 5. XRD patterns of the Cu-coated W powder and the sintered samples.
Fig. 5 shows the XRD analysis of initial powder and the sintered samples, and no impurities are observed. As shown in the inset of Fig. 5, the intensity of W peak for AN-7 is higher than the one for powder. This is because the coreeshell structure of the Cu coated W powder. But the intensity of W peaks for NM-1 is weaker than AN-7. This is because submicron particles were formed at the central region under extreme high pressure, leading to reduction of
Fig. 6. Light microscopy image of WeCu composites AN-7 (a); magnified views of the microstructure at the center (b) and the peripheral region (c) of AN-7. Light microscopy image of WeCu composites NM-1 (d); magnified views of the microstructure at the center (e) and the peripheral region (f) of NM-1.
Fig. 7. A typical image of pit due to weak bonding between W and Cu (a); magnified view of nano-sized particle formed at the center of NM-1 (b).
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Table 1 The properties of WeCu composites sintered with various conditions. No.
Shock pressure (GPa)
Sintered density %
Yield stress (MPa)
Thermal conductivity (W/mK)
AN-7 NM-1
1 12
98.9 98.4
554.52 504.64
144.82 145.62
intensity. The reason for the formation of nano-sized particles will be discussed in the section below. The W peaks from XRD profiles for initial powder, AN-7 and NM-1 were fitted using Pseudo-Voigt Function to estimate the peak widths. The results show no obvious peak broadening for the sintered samples. That implies that the submicron particles in NM-1 are formed by mechanical fragmentation, not grain refinement. 3.2. Microstructure analysis Fig. 6(a) and Fig. 6(d) are light microscopy images of AN-7 and NM-1, respectively. Fig. 6(c) and (f) show uniform distribution of W particles in Cu matrix, where no aggregations of W were observed. There are numerous dark patches distributed across the cross section of AN-7, and at the peripheral region of NM-1 as well. The samples are essentially free of porosity with a measured density > 98%T.D. (theoretical density). The magnified view by SEM shows the dark patch indicated by arrow in Fig. 7(a) is not a pore but pit. The formation of pores during shock consolidation is quite different from conventional powder metallurgy method. For shock consolidation, pores are eliminated through plastic deformation and particle fracture which are driven by shock wave instantaneously. So the pores in WeCu powders tend to retain between W particles due to insufficient deformation of W, which has been verified in hot-shock consolidation of WeCu [8]. But in this work, the WeW contact is avoided by using Cu-coated W powder. Furthermore, the size of dark patch is too large to be a pore, especially under the high pressure and high strain rate condition. However, it does exist pores due to the shrink of melted Cu under extremely high cooling rates, as indicated by the circle in Fig. 7(a). The angular pit is the evidence where a W particle existed, indicating a weak bonding between W and Cu. It is believed the weak bonding mainly comes from coating process, and the oxidation during the storage and packing stage makes it worse. The high
Fig. 8. Magnified view of center region of NM-1 shows various rupture modes: (a) fragments due to jetting penetration; (b) fragments due to impact between W particles; and (c) crushed zone at the surface.
depositing rate of Cu due to high stirring temperature lead to coarse Cu layer during the coating process, as shown in Fig. 1(b). The coarse Cu layer and the low surface energy of W particle due to its relatively large size are responsible for the weak bonding between Cu layer and W particle. The densification mechanism of Cu-coated W powders is different from the one of hot-shock consolidation with mechanically mixed WeCu powder, which is dominated by densification of W particles [8]. The densification of W particles is dominated by shock pressure; it means high-pressure is needed to promote void collapse and plastic deformation of W particles. For Cu-coated W powders, the sintering is only related to CueCu sintering, which can be easily achieved under relatively low shock pressure without preheating. This indicates that the bonding of W and Cu, which is crucial to the performance of composite, is determined by the electroless coating process. As shown in Table 1, high relative density (>98%) is obtained with a shock intensity of ~1 GPa. However, the sintering conditions required for hot-shock consolidation method [8] are high pressure (>3 GPa) and high preheating temperature (900e1000 C). By using the Cu-coated W powders, sintering mode is changed and there is no requirement of preheating. Therefore, the shock pressure is reduced and the density of the composites is improved. On cylindrical shock consolidation shock waves obviously converge and the pressures grow accordingly toward the powder center. But, during passage of the shock wave, the powders become deformed and get sintered. Energy is consumed and the shock wave should attenuate while advancing. A combination of the two effects leads to various shock front shapes, which determines the uniformity of compaction. The ideal case is when the two opposing effects counterbalance, with a hollow cone-shaped shock front as shown in Fig. 3. In such a case a uniform pressure distributes over the crosssection and a uniform consolidation is obtained. The AN-7 shows a uniform microstructure across the cross section, as shown in Fig. 6 (a), (b) and (c), indicating an ideal condition. If the amount of explosive or its detonation velocity is too high, as the case of NM-1, the Mach-stem will be formed. Under such conditions, the pressure increases continuously towards the center, cracking and melting may occur due to the extremely high pressure to tens or hundreds of GPa. This phenomenon has been observed during consolidation of metallic powders and has been studied in Refs. [19,20]. At the peripheral region for both AN-7 and NM-1, the morphologies are quite similar with the W particle in the size identical to initial powders, comparing Fig. 6(c) with Fig. 6(f). At the center region, NM-1 shows a fully densified structure without any pits or pores, as shown in Fig. 6(e). Furthermore, plenty of ultra-fined W particles are observed in Fig. 6(e), the size of which is much smaller than initial powder and is identified as nanoscale by Fig. 7(b). It is evident, from the magnified view in Fig. 8, that these nanosized W particles were formed due to the particle crushing and jet penetration. The schematic diagram in Fig. 9 shows the formation of nano-sized W particle at the center region in NM-1. While two Cu-coated W particles impact each other at high velocity driven by extremely high pressure, a Cu jet will occur due to the oblique impact of Cu layers. The ejected jet penetrates adjacent W particles at extremely high velocity, and particle crushing occurs, rather than deformation or melting due to the brittleness and high melting
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Fig. 10. Stressestrain curves from samples of quasi-static compression test.
observed in Mach-stem for metallic powders, have been avoided due to ductility of Cu and high-melting point of W, and especially due to using coating method. 3.3. Thermal and mechanical properties of WeCu composites As shown in Table 1, the thermal conductivities are 144.82 W/ mK and 145.62 W/mK for AN-7 and NM-1, respectively, which are lower than the theoretical value (237 W/mK). A continuous network of Cu, high sintering density and good interfacial contact between W and Cu are the keys to excellent thermal properties of WeCu composites [14]. As mentioned above, the weak bonding between W and Cu should be responsible for the poor thermal performance of the composites in this work. The quasi-static compression test results, shown in Fig. 10, reveal the slight difference between the two samples. These cylindrical samples have dimensions of 10 mm (length) 10 mm (diameter). No extensometer was used and the results were corrected from loadedisplacement plots by subtracting machine effects. The yield stresses are listed in Table 1. For the explosively consolidated powder mixture compact, it is suggested that the yield strength is determined by the continuous phase [21]. In this Fig. 9. Schematic diagram of the formation of nano W particles.
point of W. The small W particles are crushed into pieces while the crushing for a big particle is limited to local region. The crushed W fragments burst out and disperse into Cu matrix. This is the primary rupture mode. Simultaneously, the impact between W particles also leads to particle crushing. The crushed zone due to the normal impact is restrained to surface and the fragments due to oblique impact are propelled into Cu matrix. But these two rupture modes are limited to local region, because most kinetic energy has been deposited as deformation and jetting of Cu. The corresponding rupture modes illustrated in Fig. 9 are indicated in Fig. 8. As indicated by the arrow (a) in Fig. 8, a concave on a W particle is filled with fragments without contact with any other particles around; it is obviously a proof of jet penetration. As indicated by the arrow (b), the splashed fragments into Cu matrix show an oblique impact between W particles. And the arrow (c) in Fig. 8 shows a cursed zone due to the normal impact, with the fragments staying between W particles. The cracking and melting, which are usually
Fig. 11. Variation of hardness of the investigated samples from center to periphery.
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Fig. 12. Fracture morphology of WeCu composites after quasi-static compression test: (a) SEM image of fractured NM-1; (b) SEM image of fractured AN-7; (c) SEM image showing smooth surface of Cu due to sliding of W and (d) backscattered electron micrograph showing a W particle pulled out from the Cu matrix (dark gray for Cu and light gray for W).
work, the samples consolidated with different pressure have similar yield stress, ~500 MPa. However, the yield strength of the compact with higher shock pressure is ~50 MPa lower than the one with lower pressure, but its ductility is slightly improved. This will be discussed further below. Fig. 11 shows the variation of the hardness across the crosssection of the sample. It can be seen that the hardness of AN-7 remains consistent throughout the consolidate indicating that densification was quite uniform. But for NM-1, the graph shows a decrease from periphery towards center and a sharp drop to 107 Hv at the center. In general, the hardness increases with decreasing particle size for particulate-reinforced metal matrix composites. But the reverse behavior also has been reported [22], where the deterioration is interpreted by the inhomogeneous distribution of submicron particle. For the WeCu composites fabricated with Cucoated W powder, no W skeleton is formed during the sintering. And the indentation deformation is subjected to Cu matrix because the hardness of W particles is much higher than Cu matrix. Only if the Cu between W particles is crowded out, the W particle will be involved in resisting to the indentation. At the periphery of NM-1 where the W particles are micron-sized, Cu is crowded out during the indentation due to the weak bonding between W and Cu. And the hardness accordingly increases due to the participation of W particle. But at the center where part of the W particles are nanosized, high bonding strength is formed due to the high surface energy of nano-particles. The high bonding strength, as well as the increased WeCu interfaces, absorbs the mechanical energy that will lead to WeCu debonding and Cu extrusion; then hardness is close to the one of Cu matrix due to the lack of W participation. The high bonding strength and increasing WeCu interfaces as a result of particle fragmentation is also the reason for the improvement in bulk's ductility. The fracture type of both samples can be distinctly concluded: axial splitting fracture parallel to the loading direction. The
fractured samples were observed by scanning electron microscopy, as the general fractographs of NM-1 and AN-7 shown in Fig. 12(a) and Fig. 12(b), respectively. This failure mode is similar to the one of explosively consolidated Ni/WeAl composites under dynamic compression [21]. The discontinuously distributed W particles dispersed in the continuous Cu matrix can be considered as rigid during plastic deformation in uniaxial compression loading because its hardness significantly exceeds that of Cu. The separation is initiated at the interface between the W particle and Cu matrix, which is shown in the elemental mapping result of Fig. 13. During compression loading, the W particles move and cluster in the Cu matrix, verified by the scratch shown in Fig. 12(c). As vertical compression continues, the W particles are crowded out and move laterally, as indicated by the arrow in Fig. 12(d). Then the cracks parallel to loading direction are created due to low bonding strength between W and Cu, and propagate through the entire sample as axial splitting. The two samples consolidated with different pressure show the similar fracture mode, indicating shock pressure has no help to improve WeCu bonding strength. 4. Conclusions The following are the principal conclusions drawn from the experimental work performed on shock consolidation of Cu-coated W powder: 1. The Cu-coated W powders are prepared by chemical reduction (electroless coating). The high depositing rate of Cu during coating process and the relatively large size of W particle are responsible for the weak bonding in this work. The impurities are introduced by oxidation during the storage and packing, which also decrease the performance of WeCu composites. 2. The Cu-coated W powders were consolidated to a high relative density (exceeding 98%) by the shock wave without preheating.
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Fig. 13. Elemental mapping of the fracture surface: (a) backscattered electron micrograph of the region with a crack; (b) elemental distribution of Cu; (c) elemental distribution of W and (d) combined elemental distribution of W and Cu.
The sintering is only related to CueCu sintering due to the homogenous distribution of Cu on the W surface; and no W skeleton is formed during the sintering. Then the strength of consolidated composite is primarily determined by the Cu phase, exhibiting a lower value. The weak WeCu bonding from coating process, as well as the impurities, leads to axial splitting as a primary failure mode and inferior thermal performance. 3. The shock waves with the intensities of ~1 GPa and ~12 GPa were applied to consolidate Cu-coated W powders, respectively. The results show no difference in the bulk's performance but big difference in microstructure. The composite fabricated under 12 GPa shows an inhomogeneous microstructure along the cross section; and nano-sized W particles are formed at the center due to the particle crushing and jet penetration, as a result of Mach effect. The XRD result shows that the nano-sized particle is the result of mechanical fragmentation not grain refinement. The fragmentation of W particles decreases the hardness at the center because good bonding is formed due to the high surface energy of nano-sized particle. Acknowledgments The authors would like to express their thanks for the financial support of National Natural Science Foundation of China under grant nos. 11172043, 11402026 and 11221202. And the authors also express the appreciation to Professor C.W. Tan from school of Material Science and Technology in Beijing Institute of Technology, for the support on powder preparation. References [1] A. Ibrahim, M. Abdallah, S.F. Mostafa, A.A. Hegazy, An experimental investigation on the WeCu composites, Mater. Des. 30 (2009) 1398e1403. [2] Z.J. Zhou, J. Du, S.X. Song, Z.H. Zhong, C.C. Ge, J. Alloys Comp. 428 (2007) 146e150.
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